Transition metal hydrides may be used as neutron moderators for small-scale highly stressed nuclear power installations (NPI) at higher temperatures [7, 35]. Zirconium hydride moderator blocks of NRE should provide structural integrity during thrust regime and remain chemically inert under bimodal power regime at power fluxes of 1-1.5kW/m3 at average temperatures of 570K and temperature gradients of ~110K across the block.

Zirconium hydride parts are typically processed by reach-through hydrogen impregnation of the heated metal preforms, or by molding metal hydride pow­ders with subsequent pressure treatment [36]. The terms “reach-through impreg­nation” and “reach-through hydrogenation” mean diffusion-controlled impregnation of metal preforms up to the certain hydrogen volume fraction. Important advantage of a compacted preform hydrogenation method is the possibility of processing pore-free

material, and relative straightforwardness of the process. Among the disadvantages of this method are prolonged cycle time required for preform hydrogen impregnation (up to several weeks for large-sized performs), and the necessity to use autoclave equipment.

Forming the parts from hydride powders offers higher efficiency production of hydride parts and the capacity for composite production. The main disadvantage of the powder-based materials is typically imperfect grain boundaries that lead to poorer corrosion resistance in gaseous media. Reaction of Zr metal with hydrogen proceeds spontaneously and is accompanied by heat generation. Such spontaneous nature of reaction before hydride formation typically leads to cracking of the impregnated metal, caused by stress generation and simultaneous embrittlement. Cracking could be prevented through controlling the impregnation speed that leads to more balanced processes of stress generation and relaxation. The qualitative process and phase formation kinetics description could be obtained from the equilibrium phase diagram of the metal-hydrogen system. e. g. ZrH diagram (see Fig.4.30).

The process may be approximately divided into three stages. First stage involves formation of a в solid solution of hydrogen in zirconium (above the temperature of polymorphic transformation). The second stage involves formation of a hydride phase (at first on a surface. then in the bulk). This stage brings bulk volume changes in the materials that generate stresses possibly exceeding the ultimate strength of the material. The third stage involves compositional homogenization of the hydride phase, accompanied by the redistribution of the stresses (both in the magnitude and in the sign). Reach-through Zr impregnation by H is accompanied by volume change. For hydrides with composition close to MeH2 (e. g. for Zr hydride) this volume change may reach up to 20 %. The magnitude of the volume change influences the stress state level in the material being hydrogenated.

Based upon the data accumulated during research into hydrogen diffusion and creep [37, 38] Research Institute “Luch” has offered an optimized hydrogenation technology.

As the zirconium hydrides have insufficient strength and fracture toughness [36], there is a risk of their brittle fracture under thermal stresses generated under oper­ating conditions. This was the driving force for research into the basic mechanical parameters of zirconium hydride. The thermal stress resistance of zirconium hydride samples processed by reach-through impregnation is characterized by highly specific temperature dependence (see Table4.7).

Thermal stress resistance (TSR) of the samples is practically stable within tem­perature range of -190 to 900 °C despite the twofold strength disparity. The stability of experimentally measured TSR values is confirmed by the calculated R criterion, as the strength increase and E-modulus decrease temperature rise are compensated by significant growth of a linear expansion coefficient. The bulk structure of a sample and surface quality of hydrides after various machining treatments, influence both strength and TSR [36]. Samples after surface polishing demonstrate the largest TSR values.

The formation of hydride composites at introduction of a metal phase to 20- 40vol% reduces martensitic grain nature (Fig.4.31) and increases thermal stress resistance by two to three times, due to a local stresses relaxation, increase of a fracture toughness and the relationship a/E (Table 4.8).

Improving of bearing capacity of ceramics with a limited plasticity in a broad tem­perature range should be done taking into account operation parameters. For fuel elements of the first sections of the HRA operating in the brittle-damage temperature range, it is necessary first of all to decrease the defectiveness of ceramic using various hardening methods capable to elevate the strength. Optimizing hardening methods by modification of stressed surface state the strength of elements can be enhanced up to 100% (Table7.1).

The optimal choice of the parameters of traditional isothermal sintering of ZrC samples in different media (hydrogen, argon, and vacuum) allows achieving a high density (not less than 95%) and the bending strength about 550MPa at sintering temperatures 2,500-2,700K [2, 14].

By the beginning of work in the 1960s, information on the properties and manufacturing technology of materials for the NRE core (based on zirconium, nio­bium, uranium carbides, and zirconium hydride) was absent or inconsistent. It was known that unlike mono compound of uranium with a low melting point (2,500 K), a fuel based on solid solutions of UC-ZrC and UC-NbC carbides with nearly sto­chastic composition can provide the heating of hydrogen up to 3,000K. Therefore, investigations of solid solutions of uranium monocarbide with isomorphous, highly refractory zirconium, niobium, and monocarbides providing high melting points and compatibility of HREs with heat carriers became the most important material tech­nology direction. The prospects of the development of UC-ZrC-ZrN fuel were also outlined. The manufacturing technology of these refractory materials was based on powder metallurgy methods.

Refractory compounds are characterized by high hardness, elastic modulus, chemical stability, heat resistance and the high brittleness caused by features of interatomic interaction with mixed ionic-covalent type and low dislocation mobility [1-3]. Graphite with low values of hardness, strength and an elastic modulus possess many times higher thermal stress resistance R > 700 K in comparison with carbide compounds.

The first data on the radiation resistance of HREs (integrity, swelling, and strength) at temperatures from 1,000 to 3,100K and neutron flux intensities up to 1015- 1016cm-2 confirmed the expediency of the choice of fuel materials based in solid carbide solutions. It was decided [4, 5] for the first time in the development of a highly reliable construction in machine building to use brittle materials, which required changing construction principles and the established concepts of strength and thermal strength. A new criterion for estimating the bearing ability of thermally loaded products was introduced, which was accepted by the scientific community worldwide [6].

A. Lanin, Nuclear Rocket Engine Reactor, Springer Series in Materials Science 170, DOI: 10.1007/978-3-642-32430-7_4, © Springer-Verlag Berlin Heidelberg 2013

The choice of technologies and parameters for manufacturing materials for improving the properties of ceramics with a limited plasticity in a broad temperature range should be performed taking its operation parameters into account. For materials of

the first sections of the HRA operating in the brittle-damage temperature range, it is necessary to increase, along with the strength, the fracture toughness, which can be done first of all by decreasing the defectiveness of the materials and producing a structural state that increases the amount of energy required to produce damage. The strength in this temperature region is increased by the following methods: elimination of structural defects by optimizing the condensation and sintering regimes [15, 16], healing of defects and thermomechanical programmed control [21], modification of a stressed surface state [39], or carbon doping [17, 18, 27], or metallic band [14]. Essential increase of strength can be reached at use of nanocrystalline technologies [40, 41].

The elimination of surface defects by healing at temperatures T/Tm > 0.5 can increase the strength appreciably. The healing of radial cracks to the depth up to half a radius of the ZrC sample begins when the crack edges contact by surface diffusion. The healing is intensified with increasing the number of contacts on one length unit and is decreased with crack edge opening 8. At T/Tm= 0.6 the healing of the cracks with 8 = 1-2 ^ is completed within several hours and strength returns to the initial level (Fig.4.32). The cracks with 8 > 3 ^ and very small number of edges contacts are not healed within the same length of time. Under thermal treatment conditions the healing process has two stages.

The kinetics of the first stage (the increased intensity of the strength reduction process) is mainly defined by surface self-diffusion of Zr atoms in ZrC. After the first stage of healing, the crack is a system of isolated cavities. Its healing is performed at a lower rate (the second stage) owing to a viscous flow of the material as the surface diffusion cannot provide their volume decrease.

Another way of ceramic materials strengthening is the formation (in surface lay­ers) of compressive residual stresses preventing the appearance and spreading of cracks. The formation of residual stresses on the surface of a sample or an article is based on relaxation of thermo-elastic stresses nonuniformly distributed along the section.

Fig. 4.32 The strength reduc­tion of ZrC0.95 samples after healing of the surface cracks with the width of crack edge opening (1—I Ц, 2—1.5 Ц, 5-3-4 /Ї) at T = 2,800 K

Fig. 4.33 The bending strength change of ZrC0.95 (1, 2) and SC2O3 (5) in relation to quenching temperature (Tq) and cooling methods. 1 radia­tion cooling; 2 gaseous helium flow; 5 cooling in silicon oil

As a result of poor ductility of ceramic materials, in contrast to metals, their strengthening is performed at rather small values of criteria Bio~10-1 e. g. by blowing off a cold gas stream over the heated sample or by radiation cooling [6]. At higher Bio values the rate of thermal-elastic stress relaxation turns to be less than the rate of their increase, this leads to cracking. The temperature range of strengthening is limited in the lower domain by the temperature of the brittle — ductile transition and in the upper domain by causing the strength decrease. The increase of temperature in the range Tb-d < T < Ts rises the strengthening effectiveness. The gain in strength of strengthened ceramics is, as a rule, 20-40 % (Fig. 4.31).

A decrease of the critical volume defects, as stress concentrators, is possible by using thermal-mechanical treatment (TMT) at T/Tm > 0.6, based on stress relaxation near the concentrators [21].

The preliminary small deformation (є < 0.15%) at the low deformation rate є’ < 10-1 s-1 (Fig.4.34a) or static loading at stresses up to 0.6-0.8 a/amax (Fig. 4.34b) can increase the strength about two times. As said above there is a wide

spectrum of volume defects in materials prepared by powder metallurgy methods. It is possible to evaluate the integral structural defects by measuring the value of temporary microstresses appearing under loading in the elastic region by broadening of X-ray lines and rejecting the defective samples. The more intensive is the line broadening with growth of load, the more defective is the material. In the temper­ature region where the macro-plastic deformation becomes possible, the short — and long-term mechanical thermal-strength parameters can be optimized by almost all methods used for metals. In the temperature region where the deformation process is controlled by the motion of dislocations, the methods are used that reduce the dislocation mobility by substructural strengthening, doping of solid solutions result­ing in the formation of stronger chemical bonds in compounds, and doping with the formation of second phases. Under loading conditions, when deformation is mainly caused by grain-boundary sliding in the case of high-temperature creep, the strength can be efficiently increased by recrystallization, providing a consid­erable decrease in the length of boundaries due to the increase in the grain size. In using strengthening methods, it is necessary to take into account that struc­tural changes intended to improve mechanical characteristics at high temperatures should not impair mechanical properties at temperatures T < Tb-d. The optimal choice of the parameters of traditional isothermal sintering of ZrC samples in dif­ferent media (hydrogen. argon. and vacuum) allows achieving a high density (no less than 95%) and the bending strength about 550MPa at sintering temperatures 2,500-2,700 K [16].

NRE reactors of different Nuclear Engine Power Installations (NEPI) based on tested technologies have actually no alternatives in deep space investigations with the help of unmanned space probes and piloted interplanetary devices [1]. The developments of NREs were stopped only temporarily because humankind could not afford to spend a huge amount of money for large-scale cosmic studies. It is most likely that the NRE program should become an international collaboration program in the future, like programs for the development of the international thermonuclear experimental reactor (ITER), high-power accelerators, etc. The program would have a global character, in view of its legal. Ecological and other aspects, because the program involves the use of fission materials of the highest armament quality, it should be under international control. NRE reactors for different purposes would be based on different HRE geometries with fuel compositions providing the required efficiency. To fulfill the tasks in cosmic flights, the NRE reactor should have high reliability, with the no-failure probability no less than 0.99.

The different versions of a device operating in many regimes and capable of producing, along with the reactive thrust, the electric energy for ensuring the activity of a spacecraft were extensively developed beginning from the early 1980s [2].

They should work not only in the basic, engine regime but also in two energy regimes at low power for a few years and high power during approximately half of the specified time resource of the engine regime. The high-power regime presents no difficulties in tests because, according to all its parameters, the reactor is loaded much more weakly than in the basic regime of the IVG-1. In the low-power regime, the heat carrier flows around the HRA only outside its housing, while the heat from HREs is transferred to the housing by radiation through heat insulation. This regime is quite different from the basic regime, when the fuel is considerably burned out (down to no less than 3-5 % of the initial amount) and the chemical composition can change due to the incongruent evaporation of materials. Therefore, the efficiency of the HRA and HRE components under these conditions requires additional investigations.

It is clear that the outlook for the development and building of active cores for nuclear space energy devices will be first and foremost related to the improvement

A. Lanin, Nuclear Rocket Engine Reactor, Springer Series in Materials Science 170, 103

DOI: 10.1007/978-3-642-32430-7_8, © Springer-Verlag Berlin Heidelberg 2013

in the construction and technology of heating sections containing fuel compositions with the heat-releasing density 40MW l-1 at elevated temperatures (above 3,200K) in the engine regime and to ensuring that fission products are kept in HREs for a few years at the temperature 2,000 K in a deep vacuum or in hydrogen-containing working substances at pressures from 0.1 to 0.2atm.

The available design and technological groundwork for working of NRE reactor testify to basic possibility of creation HRA, efficient both on regimes of the reac­tive thrust, and on regimes of long work in Brighton cycle at a power production for orbiting spacecraft equipment [3]. Over the previous years, a row of thermal energy transformation concepts in electric by bimodal [4] by a fast-neutron reactor with a lithium cooling contour or with thermal emission (accordingly NEED-2 and NEED-3) have been developed (Table8.1).

Quite probably, for that the raise of working capacity of constructive elements of the NEED engineering production of nanosize materials will be developed [5, 6]. New heat insulating materials, resistant in the hydrogen environment on the basis of fibrous, porous, and multilayered carbides and nitrides of refractory metals will be created.

The Commission of the President of Russia on the Modernization and Technolog­ical Development of Russian Economics in 2009 recommended reconsidering the question of developing a spacecraft with a nuclear rocket engine. In 2010, the Gov­ernment of Russia provided the initial financial support for the development of an outline for a project of a megawatt nuclear energy device, with the possible beginning of the module construction in 2018.

Thermodynamic studies of the solid solutions of refractory compounds with uranium monocarbide were started at the RIPRA “Luch” when no data on the thermodynamic properties in the homogeneity region at high temperatures were available in the liter­ature. Only basic approaches for estimating thermodynamic properties were known. Structural and fuel materials of the reactor core are based on refractory zirconium and niobium carbides, their solid solutions with uranium, carbide compositions with carbon inclusions [4, 5, 7], and zirconium hydrides for a moderator unit (Table4.1). These materials belong to the class of so-called interstitial phases [8]. The majority of refractory compounds possess highly symmetric cubic lattice. The majority of monocarbides, mononitrides crystallizes in FCC lattice of NaCl type, nonmetallic atoms located in octahedral positions. Melting point, elastic properties and factors of thermal expansion are structurally tolerant characteristics and, in essence, they depend on bonding energy, energy of a crystal lattice necessary for division into separate ions.

High brittleness of compounds is caused by low dislocation mobility owing to an orientation of bonding, high value of forces Peierls-Nabarro and low rate mul­tiplication of dislocations. Dislocations submit to the same laws in compounds, as in metals. However in polyatomic compounds with ionic-covalent bonding, more complex geometry of dislocations’ structure is observed. Starting stresses of disloca­tions’ movement in crystals with high Peierls barrier are on one to two orders above, than in metals. The stresses demanded for the yield beginning, in turn, exceed start stresses on two orders at the moderate temperatures 0.3 Tml.

The data on the physical-mechanical properties of many refractory compounds and their compositions (some of them were used in devices for the first time) revealed the possibilities of these materials and influenced the determination of their operating conditions and estimates of their prospects [9-11]. Unlike uranium mono compounds, fuel based on solid solutions of UC-ZrC and UC-NbC carbides with a composi­tion close to the stoichiometric composition provide the heating of hydrogen up to 3,000K.

The radiation estimation of fuel elements (FA) and structural materials resistance to irradiation damage implemented during tests in reactor IVG-1 [1] and at measure­ments of properties after irradiation. The neutron flux and test temperatures in special reactor loops varied within 1012-1015n/sm2s and 450-2,000K accordingly [2, 3]. The uranium burnout in FA at propulsion mode (PM) of NRER within 1 h attained ~5 ■ 1015 and ~2 ■ 1020flss/sm3 at a power regime during 5,000h. Resource tests were carried out at heat release levels 15-35kW/sm3 during 300-4,000s. 200 FA had been tested in all.

The general regularity of materials’ behavior under irradiation may be regarded in the monography [4] and some information on interstitial phases there is in [5, 6], but the data on irradiation changes of fuel materials properties of the solid solutions with uranium of carbides and nitrides of 4-5 groups of periodic table are absent.

The most important characteristic of the radiation resistance of HREs is their dimensional stability. The swelling of HREs made of compositions based on solid solutions of uranium, zirconium, and niobium carbides depends on the fission density and irradiation temperature in a reactor is nonmonotonic function [3]. An increase in the irradiation dose of UC + ZrC + NbC and UC + ZrC fuel compositions up to 2 x 1019 fissions per cm3 at the irradiation temperature T = 1, 100 K leads to swelling by 5 % and an increase in the electric resistance up to 80 % due to the accumulation of radiation defects (mainly vacancies) (Fig. 5.1).

The bubble swelling change of fuel volume during an irradiation incubatory period is defined by balance between two competitive processes: swelling at the expense of accumulation of radiation-induced defects and irradiation sintering [3]. The swelling during the incubatory period poorly depends on the type of the investigated composi­tions and is defined, substantially, by an irradiation temperature. As concentration of the fission products is still the lowest, the dimensional sample modification is caused by accumulations of the dot defects, which evolution leads to gradual swelling growth under irradiation. Calculation shows that the maximum contribution of solid swelling at the expense of the accumulation of fission products, at B = 2-1020 fiss/sm3 makes 0.85 %., i. e., about 35 % from the total. Thus, dimensional changes of the investigated

A. Lanin, Nuclear Rocket Engine Reactor, Springer Series in Materials Science 170, DOI: 10.1007/978-3-642-32430-7_5, © Springer-Verlag Berlin Heidelberg 2013

compositions irradiated at T = 1,000 K and B from 8.4-1017 to 2-1020fiss/sm3 are defined by the swelling at the expense of the accumulations of radiation-induced defects (predominantly vacancies).

Gas bubble swelling of the fuel compositions UC + ZrC + NbC, UC + ZrC at temperature of irradiation T = 2,100 K, begins at a burnout A = 2-1019fiss/cm3. The volume modification AV/V0 from initial sample porosity p0 irradiated in the limits from 7-1017 to 1.8-1019 is defined by the constitutive equation:

AV/V0 = — C0P0[1 — exp(-B/B0)],

where C0, B—the constants depending basically from temperature irradiation and type of a fuel composition, P0 = initial porosity, B is a burnout.

The rise of irradiation temperature of the same fuel compositions from 1,000 to 2,100 K at constant fission density 2-1020 fiss/sm3 increases swelling to 6%, but reduces the values of lattice constant Aa/a and electrical resistance Ap/p to the initial value without irradiation [8] that is connected with sufficiently high annealing rate of radiation-induced defects (Fig. 5.2).

Noticeable changes of the fuel element electrical resistance A Р/Р are observed in the central HGA with temperature variation of the fuel elements (Fig.5.3).

The microscopic structure is changed simultaneously with swelling. The growing pores are arranged mainly on the grain boundaries and graphite inclusions. The most increment of porosity about 15-20%, is revealed for 3-fold solid solution UC-ZrC — NbC. Origination of pores increases the specific electrical resistance and decreases modulus of elasticity that proves to be true to the Р and E calculations on porosity at known relationships [2].

The research of the carbide fuels, implemented on the basis of the developed express complex techniques [1], confirms the radial fuel inhomogeneity swelling effect on a magnitude and on a stress sign of the fuel surfaces as at high-and low-temperature sections. Dependence of lattice constant change a0 = f(R) along fuel radius is close to the parabolic law [8, 9] (measurements of a0 were car­ried out by X-ray survey on various sections of a strongly oblique fuel section) (Fig. 5.4).

Different concentrations of radiation defects along the HRE radius (higher on the surface and lower at the more heated center) due to annihilation of point defects lead to a parabolic change in the lattice space. This gives rise to compressing stresses in the surface layers resulting in an increase of the strength and thermal strength. UC-ZrC and UC-ZrC-NbC HREs irradiated by the neutron flux J = 1014 — 1015/cm2/s at the fission density B = 1016 — 1018/cm3 and T < 0.4Tm enhances the strength by 30-50% and the thermal strength resistance by 70-80%. In this case, the strength increment decreases upon increasing the fission density above 1018 fissions per cm3 (Fig.5.5) due to the formation of vacancy and gas pores and some amorphization of materials.

The strength loss of carbide graphite (UC-ZrC+5 mas. %) occurs at lower burnout in comparison with double and three-fold solutions. The double solution UC-ZrC at irradiation is stable to fission density not less than 2-1019 fiss/cm3.

Dependence of residual stresses ars acting on a surface at inhomogeneity swelling [8]: can be easily computed by the change of lattice constant a0 on a fuel surface and measured modulus elasticity value E

ars = [A/ (1 — v)] x [(^V/V)] = [E/(1 — v)] x [Ц — 50)]

The isochronous annealing of heated HREs at 1,500 K removes inhomogeneity of swelling and the residual stresses, returning the strength to its initial level before irradiation, whereas the same irradiation doses in the case of structural carbides cause only weak changes in the elastic modulus and strength, along with a noticeable increase in the lattice pitch and electric resistance (Fig.5.6).

Irradiation of structural carbide materials under conditions similar to those of irra­diation of fuel materials, while retaining the volume and shape of the samples, leads

to an appreciable growth of the lattice parameter and growing electrical resistance, accompanied by a minor change of strength and Young modulus [10].

Under the parameters being studied, the Zr carbide density increased, while that of Nb carbide shows virtually no significant change. Lattice parameter and electrical resistance of ZrC increase much more significantly that those of NbC, while micro­hardness of these materials shows inverse behavior. After irradiation, the cracking resistance (as measured by change of the load P on the hardness tester’ indenter) of ZrC drops much more than that of NbC, while their thermal stability grows approx­imately by the same factor 1.7. Equimolar solid solution of ZrC and NbC behaves similarly to ZrC. though p behavior is rather close to that of NbC, while thermal stability demonstrating its own, unique behavior.

It is important to note that an absolute gain of electrical resistance after irradiation is much higher in ZrC and NbC, as compared to that in irradiated metals, which confirms the significant effect of the carbon sublattice on formation of the irradiation- induced defects as measured by electrical resistance change. Difference in ZrC and NbC behavior under irradiation should be attributable to the general differences in physical and chemical properties of IV and V group carbides, due to the electronic structure features [10] (Table 5.1).

It is also typical that, with increasing irradiation temperature to 1,300 K, the dif­ference in ZrC and NbC behavior remains virtually the same as at 450 K, though radiation damage consequences become much smaller. Particularly, noteworthy is also the fact of increased thermal stress resistance, possibly attributable to gen­eration of residual compressive stresses and to the effect of radiation-induced healing.

Radiation-induced fuel rod healing phenomenon has been experimentally con­firmed by declining effect of the introduced surface cracks on UC-ZrC and UC-ZrC — NbC fuel rods after low-power irradiation inside the He-filled ampoules. The degree of crack healing was estimated by monitoring changes in strength a. Electrical resis­tance p and an elastic flexure f of the fuel rods in the middle under consolidated mass impact.

The results of tests show that purely thermal annealing at relatively low (for carbide) temperature of 1,100 K cannot lead even to partial healing of the surface

cracks (see Table5.2). On the other hand, complete crack healing was observed after irradiation T = 1,100 K or after additional irradiation.

It is particularly apparent that crack-induced halving of fuel rod strength upon irradiation was replaced by 98 % strength rebound, of which 38 % correspond to a irradiation reinforcement observed in crack-free fuel rods, and the balance 60 % are attributable to an additional structural defect healing in rod surface, due to some of the critical surface defects being consumed during crack generation. The observed accelerated defect healing under irradiation is probably attributable to ‘displacement

peaks’ having much smaller size than the thermal cracks. The strength gain did not disappear after additional annealing at temperature above T^, whereas the flexure and electrical resistance returned to their initial values (i. e., before irradiation and crack introduction).

It should also be noted that the results of irradiation researches on the stability of the thermoelectric temperature transducers, for the determination of the possible temperature measurement errors. It was shown [11] that reversible and irreversible changes of the thermal electromotive forces are manifested under act of reactor irradi­ations. The reversible changes caused by additional energy release in small volume of a thermojunction are negligible small, irreversible changes grow appreciably with an increase of fluence neutrons and errors components should be considered necessar­ily as the regular capacity. The regular component of an error measurement ATr for thermocouple tungsten-rhenium VR5/20 is defined by an aspect of dependence [11]:

ATr = av фт +a2Taeaj ■ Tr,

where фТ, Фь—fluences of thermal and fast neutrons and Tirr is an irradiation temperature.

Materials for NREs should have, along with high melting points, low evaporation rates and should weakly interact with hydrogen. Changes in the material composition caused by evaporation or interaction with hydrogen should not remove the compo­sition from the homogeneity region during a specified operation time resource.

Melting temperature of carbides in the homogeneity range does not change monotonously. Typically, the maximum melting temperature is demonstrated by the non-stoichiometric phases with composition close to C/Me ratio of 0.8. The melting temperature of solid monocarbide solutions is expected to change similarly in the

homogeneity range [12]. The melting temperatures of fuel compositions based on ZrC, NbC, and ZrN decrease as the UC content increases (Figs. 4.1 and 4.2). The measured absolute values of melting temperatures for nuclear fuel rods are close to solidus line for the ZrC-UC and ZrC-NbC-UC solid solutions and average around 3,570 and 3,520 K, respectively, being significantly higher than carbon nitride melting temperature. Eutectic compounds from the typically used range of carbide-carbon compounds have the following melting temperatures: ZrC + C: 3,180; NbC + C: 3,580; TaC + C: 3,715K; the said melting temperatures being lower than that of pure carbides. Evaporation processes have a considerable effect on the performance of the HGA parts and on the behavior of these materials at high temperatures under processes involving mass transfer through vapor phase. In considering evaporation rates of alloys and compounds one should take into account not only the integrated evaporation rates, but also the partial evaporation rates. This is particularly impor­tant for such compounds as carbides, nitrides, etc., that demonstrate broad scatter of partial component evaporation rates [13]. The methods for determining the par­tial thermodynamic functions of two — and three-component interstitial phases, based on statistic and thermodynamic approach are developed in [2]. These methods can determine changes of gas pressure in the homogeneity range of congruently evaporat­ing phase compositions (i. e. the compositions which remain practically unchanged during evaporation) (Table4.2).

Predicted values for NbC are in good agreement with the experimental data; for ZrC one has to take into account the concentration dependence of the vacancy formation energy [2, 13].

As uranium has higher thermodynamic activity in the U1-y MeyC solid carbide solutions, compared to that of Zr, Nb, or Ta, and as there are no congruently evapo­rating compositions in ternary systems, consequently, the uranium loss is expected to prevail in evaporating from an open surface, with corresponding surface enrichment in Zr, Nb, or Ta atoms. Simultaneously, the condensate deposited on a cold wall should have higher U/Me ratio than the one in the initial carbide.

Noncongruent nature of U1-y MeyC solid solution evaporation makes rather prob­lematic their prolonged use at high temperatures in vacuum. Special protection mea­sures are required. In vacuum metallic component of NbCx evaporates much slower than that of zirconium carbide and, despite larger cross-section values of thermal neutrons absorption in Nb (as compared to Zr), NbC may be effectively used in carbide-graphite compositions, in binary or ternary carbide solutions, and most par­ticularly, as protective coatings on graphite parts.

As the graphitic materials are used in the HIP, irradiation durability of some graphite was investigated, especially at high temperatures 1,400-1,700 K [12]. First of all, the shrinkage rate of graphite is the most interest. For many types of graphite (except the isotropic MPG-6 graphite based on noncalcined coke) at high irradiation tem­peratures from 1,400 to 1,700K in the interval of the neutron fluxes under study (up to 2.75-1025 m-2), the shrinkage of samples was observed in both the parallel and perpendicular directions of the cut with respect to the formation axis. A decrease in the irradiation temperature to 1,200-1,500K reduces the shrinkage rate and the anisotropy in the change of geometrical dimensions (Fig. 5.7).

At the same time, the relative change in the specific electric resistance for all graphite decreases, while the relative change in the elastic modulus increases. Such a behavior of graphite materials can be explained by considerable damage to the microscopic structure and the appearance of pores and cracks, indicating ’secondary swelling.

The statistical thermodynamic theory of three-component interstitial phases devel­oped at the institute and the established properties of diffusion-controlled processes [12] such as sintering, nitration, carbonizing treatment, and oxidation facilitated the optimization of technologies used in investigations. Diffusion processes have a significant impact on the HGA materials’ interaction with hydrogen, and on consol­idation of carbide powders [3]. Experimental diffusion data were used to estimate the load-bearing capacity of fuel rods at high temperatures and the healing of surface defects. Kinetics of diffusion phenomena in the majority of cases is controlled by a number of elementary processes, whose individual contribution is rather difficult to differentiate.

Thus, densification rate during sintering and creep rate are controlled by volume, boundary, and surface diffusion [12], and by dislocation sliding. Two-component or multicomponent nature of interstitial phases generally complicates the analysis of diffusion-controlled processes, as the contribution of metallic and nonmetallic atoms in various temperature ranges may vary. Lastly, the comparison of the experimental and the theoretical data on diffusion-controlled processes requires clarification of the concept of an effective self-diffusion factor Def which is usually introduced for the multicomponent bodies [2].

For a single-component body D is a factor of a self-diffusion; for two — or multi­component bodies one should use substitute Def into the respective equations:

Def = (D*aD*b/CbD*a + CaD*b) ■ g,

where g is the thermodynamic factor, Ci is concentration of a specific component, Di is partial (chemical) self-diffusion factors that are typically defined through exper­iments with radioactive isotopes. If the diffusion mobility of one component in the alloy or compound is non-negligible (for example, diffusion mobility of interstitial elements), and D*b > D*aD, then the equation for the Def may be simplified as follows:

Def — (1/Ca) ■ D*Ag

In this case the rate of diffusion-controlled processes is defined by diffusion rate of the slowest diffusing element [2]. The developed statistic and thermodynamic the­ory of three-component interstitial phases, and our research of diffusion-controlled processes (creep, sintering, nitriding, carburization, and oxidation) make it possible to formulate evaluation recommendations concerning composition, processing, and operating conditions of NRE fuel elements (Table4.3).